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Crystallization characteristics in heavily B2H6-doped amorphous Si thin films
Journal of Alloys and Compounds 801 (2019) 352e359
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Crystallization characteristics in heavily B2H6-doped amorphous Si thin films Ji-Su Ahn a, Seung-Ki Joo b, Ashkan Vakilipour Takaloo b, Soonduck Kim c, Deok-kee Kim a, * a b c
Electrical Engineering, Sejong University, 209 Neungdong-ro, Gwangjin-gu, Seoul, 05006, South Korea School of Materials Science and Engineering, Seoul National University, Gwanak-ro, Gwanak-gu, Seoul, 08826, South Korea Department of Preventive Medicine, College of Medicine, Korea University, Seoul, 02841, South Korea
a r t i c l e i n f o
a b s t r a c t
Article history: Received 22 February 2019 Received in revised form 3 June 2019 Accepted 8 June 2019 Available online 12 June 2019
We have investigated the crystallization in amorphous silicon doped with B2H6 for varing doping levels under 3-field. Crystallization rate is classified into 3 stages based on the doping time, the mechanisms involved, and the rate determining step. The crystallization rate is saturated at around 10 mm/h after 2 min doping time. The main mechanism for the rate saturation is due to the high internal 3-field caused by the heavy doping. The direction of the internal 3-field in the case of p-type doping is opposite to that of n-type doping, which causes the final saturated crystallization rate to be higher in the case of p-type doping, and lower in the case of n-type doping, compared to the intrinsic a-Si. The B2H6-doped samples with an 3-field show an uni-directional needle network microstructure with a preferred orientation of <111> direction in (110) plane. © 2019 Elsevier B.V. All rights reserved.
Keywords: Polycrystalline silicon Microstructure Crystallization rate Boron doping Doping time
1. Introduction Fabrication of polycrystalline silicon (Poly-Si) thin film on glass substrate has been attracting a great deal of interest due to its widespread application in large-area electronics such as solar cell, high-resolution liquid crystal display, and light emitting diodes [1e4]. Recently, they have also been applied to wearable display and health care device sensors [5]. To date, Poly-Si with large crystallite grains has been achieved through a variety of techniques such as: excimer laser annealing (ELC) [6], solid phase crystallization (SPC) and rapid thermal annealing (RTA). However, these techniques have the drawbacks of high costs, high temperature processing, and non-uniform small grain size [7e10]. Metal-induced crystallization (MIC) is known as an inexpensive technique for lowering the thermal budget of amorphous silicon (a-Si) crystallization using catalytic metals such as nickel, silver, aluminum, gold, palladium, etc. in comparison to SPC [11e14]. Among the various catalytic metals used for metal-induced
* Corresponding author. E-mail address: [email protected] (D.-k. Kim). https://doi.org/10.1016/j.jallcom.2019.06.105 0925-8388/© 2019 Elsevier B.V. All rights reserved.
crystallization (MIC) [15e18], nickel (Ni) has been the preferred choice for the fabrication of polycrystalline thin film silicon [19]. However, the formation of Ni contaminations such as Ni silicide has been shown as the main drawback of the Ni induced crystallization technique which leads to the degradation of electrical performance of device [20]. To suppress the effect of Ni contaminations, metal induced lateral crystallization (MILC) has been developed where, at the edges of the Ni covered region, NiSi2 nodules move laterally and any a-Si along the path of the moving nodules will be crystallized [21e23]. So far, numerous scientific studies have been carried out in order to investigate the development of n- and p-type Poly-Si through MILC technique for the fabrication of high performance thin film solar cells and thin film transistors (TFT) [23e30]. Moreover, several scientific studies have been conducted to assess the effect of different parameters such as dopant species, applied electric field (3-field) and annealing temperature on the saturation rate and crystallite uniformity of both n- and p type-doped Poly-Si using Ni MILC [31e36]. It is reported that the combined effect of the dopants and the electric field on MILC growth is technologically important since it is relevant to TFT fabrication, specifically such procedures as the longoff-set TFT process [37]. Therefore, in our previous work [34], we systematically studied the effect of different levels of PH3-dopant
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concentrations under applied 3-field on the saturation growth rate and morphology of Poly-Si thin films during the Ni induced lateral crystallization process. However, up to now, there has been no investigation into the effect of different B2H6-dopant concentrations under applied 3-field on the saturation rate and quality of crystalline Si thin film as a result of the Ni induced lateral crystallization process. In this study, we report the Ni induced lateral crystallization of diborane (B2H6) -doped a-Si under an externally applied 3-field. Effect of doping time and electric field direction on the saturation growth rate of Ni induced crystallization of p-type doped a-Si is compared with our previously reported n-type doped a-Si. Additionally, microstructural changes of the MILC region was examined as a function of B2H6 doping time and the external 3-field direction and the saturation rate of MILC growth for B2H6-doped a-Si thin films as a function of dopant time was explained in view of the internal field and charge vacancy migration mechanisms [38].
2. Experimental section Low-pressure chemical vapor deposition (LPCVD) was used to prepare a-Si thin films with the thickness of 50 nm using Si2H6 gas, following a similar process as our previous report [24]. The deposited a-Si thin films were doped with B2H6 (p-type). The doping level in the a-Si thin films was controlled by varying times (0e5 min) of the ion mass doping (IMD). The diameter of the chamber used for IMD doping was around 30 cm with the diameter of the active region of 15 cm. The current measurements in Fig. 1(a) indicates the different p-type doping levels depending on different IMD doping times. 200-nm-thick Pt cathode and anode electrodes were deposited by the sputtering and lift-off process. 5-nm-thick Ni island patterns with an area of 200 mm 200 mm and a pitch of 715 mm were
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fabricated by sputtering and lift-off as well. Fig. 1(b) shows the schematic of a sample with an area of 1 cm 2 cm that was used in this study. Annealing at 550 C was performed for 5 h using a vacuum tube furnace. A 100 V/cm 3-field was applied to enhance the crystallization during annealing. A power source (Keithley 237) was used to apply the 3-field uniformly between the anode and the cathode, and a multi-meter (Keithley 2000) was used to measure currents during the annealing of the a-Si thin films using the furnace. Optical microscopy (OM) and field-emission scanning electron microscopy (FEeSEM) was used to analyze the microstructure of the samples. Secco-solution was used to etch the samples prior to the FE-SEM observation to enhance the a-Si and crystalline silicon (c-Si) interface and the grain boundaries in c-Si.
3. Results and discussion Fig. 1(a) shows the current measurements between the anode and the cathode in the B2H6-doped a-Si thin film (schematically shown in Fig. 1(b)), while the doping time was varied from 0 to 5 min. p-type samples with different concentrations of dopants were subjected to a constant thermal annealing at 550 C for 5 h with the 100 V/cm 3-field applied during the current measurement. As it is seen in Fig. 1(a), for the intrinsic sample, the current was not changed during the 5 h annealing at 550 C. However, for the cases of 1 and 2 min of doping, the current raised up slightly as the annealing time increased, while for the cases of 4 and 5 min of doping, which is considered a relatively high level of doping, a significant increase in current was observed. This means that the magnitude of the conductivity as well as the dependence of the conductivity is different for various doping levels. For low doping rate, annealing time did not significantly increase the current flow in Si film due to a limited amount of thermally activated carriers
Fig. 1. (a) Current through the B2H6-doped samples for varying times while annealing at 550 C for 5 h with an 3-field of 100 V/cm. (b) Schematic diagram of the a-Si sample used for the study (c) Initial and final current measurements for the IMD doping time from 0 to 5 min. (d) Initial and final resistivity measurements for the IMD doping time from 0 to 5 min.
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(B). However, for high doping rate, a remarkable increase in current flow was observed due to the involvement of exponentially increased quantity of thermally activated carriers (B). The trend of current vs. doping time for B2H6-doped samples in Fig. 1(a) was similar to that for PH3-doped samples, which was reported previously [34]. Fig. 1(c) shows the current gap between the initial and final steps of annealing treatment in different B2H6-doping times. As the IMD doping level increased, the difference between the initial and the final current flow was increased, which is attributed to the increased number of thermally activated carriers as well as MIC and MILC growth. Compared to those of the PH3-doped samples [34], the enhancement of initial and final current flow for p-type samples annealed at 550 C for 4 or 5 min doping time was less. Fig. 1(d) shows the initial and the final resistivity values corresponding to the current measurements in Fig. 1(c). As shown in Fig. 1(d), the resistivity of p-type Si thin film declined as the doping time was raised. This phenomena can be caused by the scattering of drifting electrons with boron atoms. Fig. 2 shows optical microscope (OM) images of MIC, MILC and a-Si regions of the a-Si samples subjected to different B2H6-doping times for (a) 0, (b, c) 1, (d) 2, (e) 4 and (f) 5 min followed by annealing at 550 C for 5 h. Fig. 2(b) belongs to the sample annealed without applying any 3-field, while the other ones belong to the samples with the applied 3-field of 100 V/cm. OM images show that the MIC occurred on the regions that had been selectively covered by deposited nickel thin film, whereas the uncovered area surrounding the Ni thin film was subjected to MILC phenomena for several tens of micrometers. The MILC growth rate in Fig. 2 was determined by dividing the final advancement distance by the annealing time. OM images showed that the areas of MILC in B2H6-doped samples were larger compared to those of the PH3-doped samples [34], which is generally associated with the enhancement of NiSi2/c-Si interface
velocity. In the case of B2H6-doped samples, the contribution by the activation in the c-Si region to the total current is expected to be more, while, in the case of PH3-doped samples [24], that in the a-Si region to the current is expected to be more. Fig. 2(a) represents a flat boundary between the MILC region and a-Si for the undoped sample, whereas a wavy boundary is observed in the case of B2H6-doped samples regardless of the doping time and applied 3-field. This difference in shape of interface are quite similar with our previous work (PH3-doped sample) [34]. The rugged boundary between the p-type and a-Si is thought to be due to the existence of various competitive forces affecting the MILC growth rate, such as repulsive forces between charged Ni ions and activated dopants as well as internal interfacial 3-field. As seen from Fig. 2(a), the MILC growth rate was faster in the region facing the anode by 14.5 mm/h, whereas it was slower in the region facing the cathode by 5.8 mm/h. This acceleration of the MILC rate of Si in the anode direction can be explained by “Ni ion Ni vacancy hopping model” [38,39]. In this mechanism, MILC growth is illustrated at the atomic scale as consisting of four steps. In the first step, Ni atoms in NiSi2 diffuse towards the NiSi2/a-Si interface since the Gibbs free energy of a Ni atom at the c-Si/NiSi2 interface is higher than that of a Ni atom between NiSi2/a-Si interface [40]. In the second step, Si atoms adsorption occurs at the NiSi2/a-Si interface by Si hopping in the aSi region. In the third step, Ni atoms at the c-Si/NiSi2 interface migrate through NiSi2 by charged Ni ion [41], and Ni vacancy hopping. Finally, Si atoms are rearranged at the c-Si/NiSi2 interface, and crystalline Si is formed by the atomic rearrangement of Si atoms. In the case of the anode direction, the applied 3-field facilitates the migration of the charged Ni ions toward the NiSi2/a-Si interface enhancing the MILC growth rate, while, in the cathode direction, the charged Ni ions drift away from the leading edge of NiSi2 by reversing the direction of the applied 3-field, retarding the growth of MILC.
Fig. 2. OM micrographs of the B2H6-doped a-Si thin films doped for (a) 0, (b, c) 1, (d) 2, (e) 3, and (f) 4 min after the 550 C anneal for 5 h. (b) Corresponds to the sample without the 3-field, while others correspond to the case with 3-field.
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Fig. 2(b) shows the plane-view OM image of the a-Si sample after 1 min of B2H6 doping and subsequent annealing at 550 C for 5 h without any applied 3-field. The MILC rate is observed to be 10.1 mm/h in both anode and cathode directions which is almost double compared to our previous study on PH3-doped samples [39]. Based on the hopping mechanism, the repulsive and attractive driving forces between the charged Ni ion and the activated dopants (B and Pþ), the progress of the movement of c-Si/a-Si interface can be promoted or demoted. Hence, the MILC growth rate is accelerated or decelerated by p-type doping and n-typedoping, respectively. Based on Fig. 2(c), the MILC growth rate of B2H6-doped sample was found to have the opposite dependence on the applied 3-field in anode and cathode directions. In the anode direction, the MILC growth rate was observed to be 11.8 mm/h which means 16.8% enhancement compared to the case without the applied 3-field. However, in the cathode direction, the MILC growth rate increased remarkably to 14.4 mm/h (42.5% enhancement) compared to the case without the applied 3-field. Moreover, it was observed that, the electric field direction dependent of MILC growth rate of B2H6doped sample are opposite to PH3-doped samples [34]. The MILC growth rate's enhancement of B2H6-doped sample facing the cathode direction can be explained by the charge vacancy migration mechanism [38]. It is known that single vacancies in Si occur in four ionization states (Vþ, V2þ, V, V2) whose concentrations are dependent on temperature and Fermi energy level based on the following equations [42]. [Vþ] ¼ [Vo] exp ((Evþ - EF)/KT)
(1)
[V] ¼ [Vo] exp ((EF- Ev-)/KT)
(2)
[V2] ¼ [Vo] exp ((2EF - Ev- - Ev2-)/KT),
(3)
where Ev, Ev-, and Ev2- are the energy levels of Vþ, V, and V2. This change in concentration can illustrate the enhancement of the MILC rate of the B2H6-doped sample in the cathode direction. In the case of a p-type sample, concentration of dominant Vþ is increased in the cathode direction by migration of positive charge vacancies in the reverse direction of the 3-field. Consequently, the growth rate in the cathode direction is accelerated due to the promotion of the adsorption of the Si atoms at the a-Si/NiSi2 and the Si atom rearrangement at the c-Si/NiSi2 phase boundary, which was found to be opposite compared to that for the n-type sample [34,43,44]. For doping time above 2 min, the MILC growth rate decreased toward a saturated value as shown in Fig. 2(d)-2(e), and, finally, it was saturated for the 5 min doping sample. Fig. 3(a)e3(c) show the FE-SEM images of morphology and microstructure of MILC and a-Si region corresponding to the 1 min B2H6 doped sample after 5 h annealing at 550 C, with the corresponding magnified FE-SEM images in the insets. The sample annealed without any 3-field is shown in Fig. 3(a), whereas those with applying external 3-field during annealing treatment in the cathode and the anode directions are shown in Fig. 3(b)e(c), respectively. The marks in Fig. 2(b)e(c) show the locations where the FE-SEM images in Fig. 3(a)-3(c) were taken. Preferred crystal growth orientations and the microstructures of the MILC growth region of 1min B2H6-doped sample with and without the applied 3-field in both the anode and the cathode directions are summarized in Table 1. The maximum crystallite lengths measured for an area of 10 mm 10 mm are given in Table 1 as well. The microstructure of the MILC region of the B2H6-doped sample without the 3-field showed a network of bi-directional
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needle like structures of branched crystallites with a regular angle of 70 to 110 between primary and secondary branches. The orientation of these branches are expected to be related with the orientation of NiSi2 precipitates in the a-Si film determining the subsequent lateral growth <111> direction of the crystallite [11,44]. The maximum length of branched crystallite was over 6.9 mm, with the width of 40e50 nm which differs from previous study results [37]. The microstructure of the B2H6-doped sample annealed in the anode direction of the 3-field showed an uni-directional needle like structure with preferred crystallite growth orientation of the <111> direction in the (110) plane [45]. Unlike samples grown without the 3-field (Fig. 3(a)), in the case of samples grown with the 3-field, the branching secondary arm is highly restrained, and a uniform array of parallel crystallites was formed as shown in Fig. 3(b)e(c). The length (over 11 mm) and the width (50e60 nm) of the crystallites with the 3-field at the leading edge of the MILC growth in the anode direction was 36% longer and 20% thicker, respectively, than those (6.9 mm and 40e50 nm) without the 3-field. The length (over 12 mm) and the width (60e80 nm) of the crystallites with the 3-field at the leading edge of the MILC growth in the cathode direction was 42% longer and 33% thicker, respectively, than those (6.9 mm and 40e50 nm) without the 3-field. This phenomenon may be attributed to the second level of Fig. 3(d) describing the different MILC growth rate of B2H6-dope a-Si sample in anode and cathode direction of 3-field along with the sample without the applied 3-field (explained in following paragraphs). The microstructure of the MILC region of B2H6-doped sample facing the cathode (Fig. 3(b)) also showed an uni-directional needle like structure. The maximum length of crystallite was over 12 mm with enhanced directionality compared to sample without electric field. This extension of crystallite length is associated with the long range migration of NiSi2 precipitates in the direction opposite the applied 3-field compared to the case without 3-field [40]. Meanwhile, the width of crystallite in the cathode direction was 60e80 nm, which was thicker than that (40e50 nm) of crystallite without the 3-field. In the case of PH3 doping [34], the MILC growth rate was decreased by increasing the level of doping, while in the case of B2H6 doping, the MILC growth rate increased compared to the PH3 doping and intrinsic Si. In the case of PH3-doped a-Si, the activated dopant (Pþ) suppresses the movement of the c-Si/NiSi2 interface through attraction of charged Ni ion which decelerate the MILC rate. However, in the case of B2H6 doped a-Si, the activated dopant (B) enhances the movement of the c-Si/NiSi2 interface through repulsion of charged Ni ion which accelerate the MILC rate. Fig. 3(d) shows the MILC growth rate of B2H6-doped a-Si thin films after 5 h annealing at 550 C in anode (þ, Blue) and cathode (-, Red) directions under an electric field (100 V/cm) together with the MILC growth rate without an electric field (Black) classified into three different MILC growth stages depending on the doping level. The MILC growth rate of the sample in anode direction was 14.54 mm/h in an intrinsic state and decreased to 11.82 mm/h after 1 min B2H6 doping. After further doping, the MILC growth rate decreased from 10.46 mm/h at 2 min to 10.24 mm/h at 4 min, and finally became saturated around 9.88 mm/h. On the other hand, without the 3-field, the growth rate increased from 7.38 mm/h to 10.08 mm/h when the doping level increased from no doping to 1 min doping. Above 1 min doping time, the growth rate became saturated around 10 mm/h. For the sample with the 3-field in the cathode direction, the growth rate was 5.78 mm/h for no doping and increased sharply to 14.42 mm/h after 1 min B2H6 doping. Above 1 min doping, the growth rate decreased to 11.74 mm/h after 2 min doping, and, finally, was saturated around 10.6 mm/h after 5 min doping. MILC growth rate of B2H6-doped a-Si thin films saturated regardless of the applied 3-field. In Fig. 3(d), the saturated MILC
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Fig. 3. (a) FE-SEM micrograph of the 1 min B2H6-doped a-Si thin films after the 550 C anneal for 5 h (no 3-field). FE-SEM micrographs of the 1 min B2H6-doped a-Si thin films in (b) the cathode and (c) the anode directions (under 3-field). (d) B2H6-doped a-Si thin film MILC growth rate for 0e5 min doping time in the anode (blue, þ) and the cathode (red, -) directions after the 550 C anneal for 5 h. MILC growth rate for the sample without 3-field is shown as black. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Table 1 Preferred Orientation and microstructure of MILC growth region for the sample doped with B2H6 for 1 min. Case
Dominant Microstructure
Expected Preferred Orientation
Maximum Length of Crystallite (mm)
Width of Needle Crystallite (nm)
Without ε-field With ε-field, anode
Uni-directional parallel growth Uni-directional parallel growth (with enhanced directionality) Uni-directional parallel growth (with enhanced directionality)
<111>/(110) <111>/(110)
>6.9 >11
40e50 50e60
<111>/(110)
>12
60e80
With ε-field, cathode
growth rate (~12 mm/hr) in the cathode direction was higher than that without an 3-field (~10 mm/hr), while that (~10 mm/hr) in the anode direction was similar. The saturated MILC growth rate (~10 mm/hr) in B2H6-doped a-Si was higher than that in intrinsic a-Si without an 3-field (7.3 mm/hr), while that (2 mm/hr) in PH3-doped aSi was lower than that in intrinsic a-Si without an 3-field (7.3 mm/hr) [34]. In the first part of Fig. 3(d), named “intrinsic stage,” the MILC growth rate of B2H6-doped samples is highest in the anode direction, intermediate without electric field, and lowest in the cathode direction. This order of MILC rate can be explained using the threestep hopping mechanism [38,43,44]. In the case of the anode direction, negatively charged Ni ions drift faster toward the leading interface edge of the anode direction, causing an enhancement of MILC growth rate. While, in the case of the cathode direction, due to the opposite direction of drift, MILC growth rate is decreased. However, in the absence of an 3-field, the MILC growth rate is elevated by the repulsive force between activated dopant (B) and negatively charged Ni ions. From Level I to Level II (“MILC rate reversal stage”), as doping
time is increased the MILC growth rate of the B2H6-doped sample in the cathode direction increased, whereas the decreasing trend in MILC growth rate was observed for the B2H6-doped sample in the anode direction. This phenomenon is associated with the positively-charged vacancy dependent MILC growth rate of B2H6doped a-Si thin film. The existence of positively-charged vacancies (Vþ and V2þ) in ptype Si can be used to explain the increasing and decreasing MILC growth rate in the cathodic and anodic directions, respectively. For p-type sample, as doping time is increased, the concentration of dominant positively-charged vacancies (Vþ and V2þ) is increased in the cathode direction due to their drifting toward the leading edge of NiSi2, which accelerates the MILC growth rate by increasing the adsorption and atomic rearrangement of the Si atoms. However, as doping time is increased more, the MILC growth rate in the cathode direction is slightly decreased due to establishment of an attractive force between activated dopant (B) and positively-charged vacancies. Decreasing the MILC growth rate as a function of doping time for p-type a-Si sample in the anode direction also can be understood with the same explanation. However, in this case,
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Fig. 4. Schematic diagrams of the MILC growth mechanisms for the B2H6-doped a-Si thin films under an applied 3-field. (a), (c), and (e) correspond to the anode direction, while (b), (d) and (f) corresponds to the cathode direction ((a) and (b) for doping level I, (c) and (d) for doping level II, (e) and (f) for doping level III).
concentration of positively charged vacancies is decreased in the anode direction. MILC growth rate is decelerated, and consequently the rate of the adsorption and the atomic rearrangement of Si atoms declines [46]. As seen from the third part of Fig. 3(d), the MILC growth rate is saturated after 3 min of B2H6 doping. This saturation behavior of MILC growth rate of B2H6-doped sample may be due to the internal interfacial electric field, similar to PH3-doped sample in our previous study [34]. Due to high doping time, the crystallization region of the Si thin film can be considered as having high concentration of activated stationary ions (B). On the other hand, a high density of states of defects in the mobility gap [47], holding holes and electrons, exists in the a-Si region, which may act as either p-type or n-type depending on the doping type. These defects can emit an extra electron and thus play a role as a donor. Therefore, it can be assumed that the a-Si region appears as a material with a donor state density (nx). Due to the existence of donor and acceptor atoms, a p-n junction is thought to be formed at the c-Si/a-Si interface. The internal 3-field at the interface can be estimated using similar equations as used for the n-type doping [34]. The estimated internal interfacial 3-field for the doping level of 1016 - 1020/cm3 was 3 104 V/cm to 3 MV/cm. The much higher internal interfacial 3-field induces the MILC growth rate saturation as the IMD doping time increases, for both with or with the applied external 3-field [34].
The operating mechanism and the rate determining step of pdoped MILC growth regions for different growth rate stages I, II, and III given in Fig. 4 are summarized in Table 2 as well. Filled and empty stars in Table 2 correspond to the factors enhancing and retarding the MILC, respectively: the greater the number of stars, the stronger the effect on the MILC growth rate. Doping Level I corresponds to the case of an intrinsic Si with or without 3-field, which is shown schematically in Fig. 4(a)e(b), where the rate of MILC is determined by the Ni ion and Ni vacancy hopping step. The case of doping Level II, which is a MILC rate reversal stage with charged vacancy migration affected by an external 3-field and slight p-type doping, is shown schematically in Fig. 4(c)e(d). The operating mechanism in doping Level I is still present, but the effect becomes minimal, which is indicated with one star. The other two major operating mechanisms, which are the repulsive interaction between B and Ni and the positively charged Si vacancies migration with the external electric field, are marked with three stars. The repulsive interaction between B and Ni enhances MILC both in the anode and the cathode direction, while the interaction between the positively charged Si vacancies with the electric field retards MILC in the anode direction and enhances it in the cathode direction. Since the two mechanisms are in the same direction in the cathode direction, the MILC rate was much enhanced in the cathode direction and resulted in the MILC rate reversal in the
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Table 2 Operating mechanism and rate determining step of boron-doped MILC growth regions for different growth stages I, II, and III. Mechanism
Doping Level I (Intrinsic Stage)
Without ε-field
① Nominal MILC by Ni ion and Ni vacancy hopping model @ 550 C
Doping Level II (MILC Rate Reversal Stage)
Doping Level III (MILC Rate Saturation Stage)
① MILC rate enhancement by B- and Ni- ① Saturation by large Internal interfacial ε-field due to repulsive interaction (+++) heavy doping (Internal interfacial ε-field » externally applied ε-field) (+++) ② MILC rate enhancement by B- and Ni- repulsive interaction (++) With ε-field - Anode ① MILC rate enhancement in the anode ① MILC rate enhancement by B- and Ni- ① Saturation by large Internal interfacial ε-field due to heavy doping (Internal interfacial ε-field » externally repulsive interaction (+++) Direction direction through Ni- movement ② Retardation of atomic rearrangement applied ε-field) (+++) enhancement by ε-field (+++) ② MILC rate enhancement by B- and Ni- repulsive by positively charged Si vacancy (Vþ) interaction (++) concentration decrease in the anode ③ Retardation of atomic rearrangement by positively direction (☆☆☆) ③ MILC rate enhancement in the anode charged Si vacancy (Vþ) concentration decrease in the direction through Ni- movement anode direction (☆) enhancement by ε-field (+) ④ MILC rate enhancement in the anode direction through Ni- movement enhancement by ε-field (+) With ε-field - Cathode ① MILC rate retardation in the cathode ① MILC rate enhancement by B- and Ni- ① Saturation by large Internal interfacial ε-field due to heavy doping (Internal interfacial ε-field » externally repulsive interaction (+++) Direction direction through Ni- movement applied ε-field) (+++) ② Enhancement of atomic retardation by ε-field (☆☆☆) rearrangement by positively charged Si ② MILC rate enhancement by B- and Ni- repulsive vacancy (Vþ) concentration increase in interaction (++) ③ Enhancement of atomic rearrangement by positively the cathode direction (+++) ③ MILC rate retardation in the cathode charged Si vacancy (Vþ) concentration increase in the cathode direction (+) direction through Ni- movement ④ MILC rate retardation in the cathode direction through retardation by ε-field (☆) Ni- movement retardation by ε-field (☆) Step 2 - Ni ion and Ni vacancy hopping Rate Determining Step Step 2 - Ni ion and Ni vacancy hopping Step 3 - Atomic rearrangement by positively charged Si vacancy (Vþ) migration
cathode direction. In the doping Level II, the rate determining step is the atomic rearrangement by positively-charged Si vacancy migration. Doping Level III corresponds to a MILC rate saturation stage with over doping effect in the a-Si region, as shown schematically in Fig. 4(e)e(f). The three operating mechanisms in doping Level II are still operating, but the effect is less significant. In the case of doping Level III, a fourth operating mechanism, which is an internal interfacial 3-field due to heavy doping, becomes the dominant operating mechanism. The internal interfacial 3-field takes over the other effects, which makes the MILC rate independent of 3-field direction. The internal interfacial 3-field becomes much greater than the externally applied 3-field in doping Level III, and the rate determining step is the Ni ion and Ni vacancy hopping step. The direction of the internal 3-field in the case of p-type doping is opposite to that of n-type doping, which causes the final saturated MILC rate to be higher in the case of p-type doping, and lower in the case of n-type doping compared to the intrinsic a-Si. In the case of p-type doping, the internal 3-field caused by the space charge region in between p-type c-Si and a-Si pushes the Ni ion away towards the a-Si region, while in the case of n-type doping, the internal 3-field pulls Ni towards the c-Si region. 4. Conclusions We have studied MILC growth rate saturation of heavily boron doped a-Si thin films by applying an external 3-field. The growth rate was classified into 3 stages based on the doping time. The operating mechanism and the rate determining step of p-doped MILC growth regions for three different growth stages are suggested. It was suggested that many competing driving forces are operating for p-type a-Si MILC growth, such as the repulsive force between the Ni ion and the activated B dopants, the forces between the Ni ions and the external 3-field, the positively charged vacancy migration, and the internal interfacial 3-field. For both the
samples with and without an applied 3-field, MILC growth rate saturated at around 10 mm/h after 2 min doping time, regardless of the applied 3-field, due to the much higher internal 3-field. The direction of the internal 3-field in the case of p-type doping was opposite to that of n-type doping, which caused the final saturated crystallization rate to be higher in the case of p-type doping and lower in the case of n-type doping, compared to the intrinsic a-Si. In the case of p-type doping, the internal 3-field caused by the space charge region in the p-type c-Si and a-Si was pushing Ni away towards the a-Si region, while in the case of n-type doping, that in the n-type c-Si and a-Si was pulling Ni towards the c-Si region. Acknowledgements This study was supported by Basic research program (2016R1D1A1B01009537) and Nano Material Technology Development Program (2015M3A7B7045496) through the National Research Foundation of Korea funded by the Ministry of Science, ICT & Future Planning. This research was supported by the MOTIE (Ministry of Trade, Industry & Energy (10080581)) and the KSRC (Korea Semiconductor Research Consortium) support program for the development of the future semiconductor device. References [1] R.E.I. Schropp, M. Zeman, Amorphous and Microcrystalline Silicon Solar Cells: Modeling, Materials and Device Technology, Kluwer Academic Publishers, Boston, 1998. [2] J.H. Choi, J.H. Cheon, S.K. Kim, J. Jang, Giant-grain silicon (GGS) and its application to stable thin film transistor Displays 26 (2005) 137. [3] A.H. Mahan, S.P. Ahrenkiel, R.E.I. Schropp, H. Li, D.S. Ginley, A comparison of grain nucleation and grain growth during crystallization of HWCVD and PECVD Si:H films, Thin Solid Films 516 (2008) 529. [4] H.C. Jong, Y.K. Do, K.C. Byoung, S.S. Woo, J. Jin, Metal induced lateral crystallization of amorphous silicon through a silicon nitride cap layer, Electrochem. Solid State Lett. 6 (2003) 16. [5] C. Pang, C. Lee, K.-Y. Suh, Recent advances in flexible sensors for wearable and implantable devices, J. Appl. Polym. Sci. 130 (2013) 1429.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Crystallization characteristics in heavily B2H6-doped amorphous Si thin films Ji-Su Ahn a, Seung-Ki Joo b, Ashkan Vakilipour Takaloo b, Soonduck Kim c, Deok-kee Kim a, * a b c
Electrical Engineering, Sejong University, 209 Neungdong-ro, Gwangjin-gu, Seoul, 05006, South Korea School of Materials Science and Engineering, Seoul National University, Gwanak-ro, Gwanak-gu, Seoul, 08826, South Korea Department of Preventive Medicine, College of Medicine, Korea University, Seoul, 02841, South Korea
a r t i c l e i n f o
a b s t r a c t
Article history: Received 22 February 2019 Received in revised form 3 June 2019 Accepted 8 June 2019 Available online 12 June 2019
We have investigated the crystallization in amorphous silicon doped with B2H6 for varing doping levels under 3-field. Crystallization rate is classified into 3 stages based on the doping time, the mechanisms involved, and the rate determining step. The crystallization rate is saturated at around 10 mm/h after 2 min doping time. The main mechanism for the rate saturation is due to the high internal 3-field caused by the heavy doping. The direction of the internal 3-field in the case of p-type doping is opposite to that of n-type doping, which causes the final saturated crystallization rate to be higher in the case of p-type doping, and lower in the case of n-type doping, compared to the intrinsic a-Si. The B2H6-doped samples with an 3-field show an uni-directional needle network microstructure with a preferred orientation of <111> direction in (110) plane. © 2019 Elsevier B.V. All rights reserved.
Keywords: Polycrystalline silicon Microstructure Crystallization rate Boron doping Doping time
1. Introduction Fabrication of polycrystalline silicon (Poly-Si) thin film on glass substrate has been attracting a great deal of interest due to its widespread application in large-area electronics such as solar cell, high-resolution liquid crystal display, and light emitting diodes [1e4]. Recently, they have also been applied to wearable display and health care device sensors [5]. To date, Poly-Si with large crystallite grains has been achieved through a variety of techniques such as: excimer laser annealing (ELC) [6], solid phase crystallization (SPC) and rapid thermal annealing (RTA). However, these techniques have the drawbacks of high costs, high temperature processing, and non-uniform small grain size [7e10]. Metal-induced crystallization (MIC) is known as an inexpensive technique for lowering the thermal budget of amorphous silicon (a-Si) crystallization using catalytic metals such as nickel, silver, aluminum, gold, palladium, etc. in comparison to SPC [11e14]. Among the various catalytic metals used for metal-induced
* Corresponding author. E-mail address: [email protected] (D.-k. Kim). https://doi.org/10.1016/j.jallcom.2019.06.105 0925-8388/© 2019 Elsevier B.V. All rights reserved.
crystallization (MIC) [15e18], nickel (Ni) has been the preferred choice for the fabrication of polycrystalline thin film silicon [19]. However, the formation of Ni contaminations such as Ni silicide has been shown as the main drawback of the Ni induced crystallization technique which leads to the degradation of electrical performance of device [20]. To suppress the effect of Ni contaminations, metal induced lateral crystallization (MILC) has been developed where, at the edges of the Ni covered region, NiSi2 nodules move laterally and any a-Si along the path of the moving nodules will be crystallized [21e23]. So far, numerous scientific studies have been carried out in order to investigate the development of n- and p-type Poly-Si through MILC technique for the fabrication of high performance thin film solar cells and thin film transistors (TFT) [23e30]. Moreover, several scientific studies have been conducted to assess the effect of different parameters such as dopant species, applied electric field (3-field) and annealing temperature on the saturation rate and crystallite uniformity of both n- and p type-doped Poly-Si using Ni MILC [31e36]. It is reported that the combined effect of the dopants and the electric field on MILC growth is technologically important since it is relevant to TFT fabrication, specifically such procedures as the longoff-set TFT process [37]. Therefore, in our previous work [34], we systematically studied the effect of different levels of PH3-dopant
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concentrations under applied 3-field on the saturation growth rate and morphology of Poly-Si thin films during the Ni induced lateral crystallization process. However, up to now, there has been no investigation into the effect of different B2H6-dopant concentrations under applied 3-field on the saturation rate and quality of crystalline Si thin film as a result of the Ni induced lateral crystallization process. In this study, we report the Ni induced lateral crystallization of diborane (B2H6) -doped a-Si under an externally applied 3-field. Effect of doping time and electric field direction on the saturation growth rate of Ni induced crystallization of p-type doped a-Si is compared with our previously reported n-type doped a-Si. Additionally, microstructural changes of the MILC region was examined as a function of B2H6 doping time and the external 3-field direction and the saturation rate of MILC growth for B2H6-doped a-Si thin films as a function of dopant time was explained in view of the internal field and charge vacancy migration mechanisms [38].
2. Experimental section Low-pressure chemical vapor deposition (LPCVD) was used to prepare a-Si thin films with the thickness of 50 nm using Si2H6 gas, following a similar process as our previous report [24]. The deposited a-Si thin films were doped with B2H6 (p-type). The doping level in the a-Si thin films was controlled by varying times (0e5 min) of the ion mass doping (IMD). The diameter of the chamber used for IMD doping was around 30 cm with the diameter of the active region of 15 cm. The current measurements in Fig. 1(a) indicates the different p-type doping levels depending on different IMD doping times. 200-nm-thick Pt cathode and anode electrodes were deposited by the sputtering and lift-off process. 5-nm-thick Ni island patterns with an area of 200 mm 200 mm and a pitch of 715 mm were
353
fabricated by sputtering and lift-off as well. Fig. 1(b) shows the schematic of a sample with an area of 1 cm 2 cm that was used in this study. Annealing at 550 C was performed for 5 h using a vacuum tube furnace. A 100 V/cm 3-field was applied to enhance the crystallization during annealing. A power source (Keithley 237) was used to apply the 3-field uniformly between the anode and the cathode, and a multi-meter (Keithley 2000) was used to measure currents during the annealing of the a-Si thin films using the furnace. Optical microscopy (OM) and field-emission scanning electron microscopy (FEeSEM) was used to analyze the microstructure of the samples. Secco-solution was used to etch the samples prior to the FE-SEM observation to enhance the a-Si and crystalline silicon (c-Si) interface and the grain boundaries in c-Si.
3. Results and discussion Fig. 1(a) shows the current measurements between the anode and the cathode in the B2H6-doped a-Si thin film (schematically shown in Fig. 1(b)), while the doping time was varied from 0 to 5 min. p-type samples with different concentrations of dopants were subjected to a constant thermal annealing at 550 C for 5 h with the 100 V/cm 3-field applied during the current measurement. As it is seen in Fig. 1(a), for the intrinsic sample, the current was not changed during the 5 h annealing at 550 C. However, for the cases of 1 and 2 min of doping, the current raised up slightly as the annealing time increased, while for the cases of 4 and 5 min of doping, which is considered a relatively high level of doping, a significant increase in current was observed. This means that the magnitude of the conductivity as well as the dependence of the conductivity is different for various doping levels. For low doping rate, annealing time did not significantly increase the current flow in Si film due to a limited amount of thermally activated carriers
Fig. 1. (a) Current through the B2H6-doped samples for varying times while annealing at 550 C for 5 h with an 3-field of 100 V/cm. (b) Schematic diagram of the a-Si sample used for the study (c) Initial and final current measurements for the IMD doping time from 0 to 5 min. (d) Initial and final resistivity measurements for the IMD doping time from 0 to 5 min.
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(B). However, for high doping rate, a remarkable increase in current flow was observed due to the involvement of exponentially increased quantity of thermally activated carriers (B). The trend of current vs. doping time for B2H6-doped samples in Fig. 1(a) was similar to that for PH3-doped samples, which was reported previously [34]. Fig. 1(c) shows the current gap between the initial and final steps of annealing treatment in different B2H6-doping times. As the IMD doping level increased, the difference between the initial and the final current flow was increased, which is attributed to the increased number of thermally activated carriers as well as MIC and MILC growth. Compared to those of the PH3-doped samples [34], the enhancement of initial and final current flow for p-type samples annealed at 550 C for 4 or 5 min doping time was less. Fig. 1(d) shows the initial and the final resistivity values corresponding to the current measurements in Fig. 1(c). As shown in Fig. 1(d), the resistivity of p-type Si thin film declined as the doping time was raised. This phenomena can be caused by the scattering of drifting electrons with boron atoms. Fig. 2 shows optical microscope (OM) images of MIC, MILC and a-Si regions of the a-Si samples subjected to different B2H6-doping times for (a) 0, (b, c) 1, (d) 2, (e) 4 and (f) 5 min followed by annealing at 550 C for 5 h. Fig. 2(b) belongs to the sample annealed without applying any 3-field, while the other ones belong to the samples with the applied 3-field of 100 V/cm. OM images show that the MIC occurred on the regions that had been selectively covered by deposited nickel thin film, whereas the uncovered area surrounding the Ni thin film was subjected to MILC phenomena for several tens of micrometers. The MILC growth rate in Fig. 2 was determined by dividing the final advancement distance by the annealing time. OM images showed that the areas of MILC in B2H6-doped samples were larger compared to those of the PH3-doped samples [34], which is generally associated with the enhancement of NiSi2/c-Si interface
velocity. In the case of B2H6-doped samples, the contribution by the activation in the c-Si region to the total current is expected to be more, while, in the case of PH3-doped samples [24], that in the a-Si region to the current is expected to be more. Fig. 2(a) represents a flat boundary between the MILC region and a-Si for the undoped sample, whereas a wavy boundary is observed in the case of B2H6-doped samples regardless of the doping time and applied 3-field. This difference in shape of interface are quite similar with our previous work (PH3-doped sample) [34]. The rugged boundary between the p-type and a-Si is thought to be due to the existence of various competitive forces affecting the MILC growth rate, such as repulsive forces between charged Ni ions and activated dopants as well as internal interfacial 3-field. As seen from Fig. 2(a), the MILC growth rate was faster in the region facing the anode by 14.5 mm/h, whereas it was slower in the region facing the cathode by 5.8 mm/h. This acceleration of the MILC rate of Si in the anode direction can be explained by “Ni ion Ni vacancy hopping model” [38,39]. In this mechanism, MILC growth is illustrated at the atomic scale as consisting of four steps. In the first step, Ni atoms in NiSi2 diffuse towards the NiSi2/a-Si interface since the Gibbs free energy of a Ni atom at the c-Si/NiSi2 interface is higher than that of a Ni atom between NiSi2/a-Si interface [40]. In the second step, Si atoms adsorption occurs at the NiSi2/a-Si interface by Si hopping in the aSi region. In the third step, Ni atoms at the c-Si/NiSi2 interface migrate through NiSi2 by charged Ni ion [41], and Ni vacancy hopping. Finally, Si atoms are rearranged at the c-Si/NiSi2 interface, and crystalline Si is formed by the atomic rearrangement of Si atoms. In the case of the anode direction, the applied 3-field facilitates the migration of the charged Ni ions toward the NiSi2/a-Si interface enhancing the MILC growth rate, while, in the cathode direction, the charged Ni ions drift away from the leading edge of NiSi2 by reversing the direction of the applied 3-field, retarding the growth of MILC.
Fig. 2. OM micrographs of the B2H6-doped a-Si thin films doped for (a) 0, (b, c) 1, (d) 2, (e) 3, and (f) 4 min after the 550 C anneal for 5 h. (b) Corresponds to the sample without the 3-field, while others correspond to the case with 3-field.
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Fig. 2(b) shows the plane-view OM image of the a-Si sample after 1 min of B2H6 doping and subsequent annealing at 550 C for 5 h without any applied 3-field. The MILC rate is observed to be 10.1 mm/h in both anode and cathode directions which is almost double compared to our previous study on PH3-doped samples [39]. Based on the hopping mechanism, the repulsive and attractive driving forces between the charged Ni ion and the activated dopants (B and Pþ), the progress of the movement of c-Si/a-Si interface can be promoted or demoted. Hence, the MILC growth rate is accelerated or decelerated by p-type doping and n-typedoping, respectively. Based on Fig. 2(c), the MILC growth rate of B2H6-doped sample was found to have the opposite dependence on the applied 3-field in anode and cathode directions. In the anode direction, the MILC growth rate was observed to be 11.8 mm/h which means 16.8% enhancement compared to the case without the applied 3-field. However, in the cathode direction, the MILC growth rate increased remarkably to 14.4 mm/h (42.5% enhancement) compared to the case without the applied 3-field. Moreover, it was observed that, the electric field direction dependent of MILC growth rate of B2H6doped sample are opposite to PH3-doped samples [34]. The MILC growth rate's enhancement of B2H6-doped sample facing the cathode direction can be explained by the charge vacancy migration mechanism [38]. It is known that single vacancies in Si occur in four ionization states (Vþ, V2þ, V, V2) whose concentrations are dependent on temperature and Fermi energy level based on the following equations [42]. [Vþ] ¼ [Vo] exp ((Evþ - EF)/KT)
(1)
[V] ¼ [Vo] exp ((EF- Ev-)/KT)
(2)
[V2] ¼ [Vo] exp ((2EF - Ev- - Ev2-)/KT),
(3)
where Ev, Ev-, and Ev2- are the energy levels of Vþ, V, and V2. This change in concentration can illustrate the enhancement of the MILC rate of the B2H6-doped sample in the cathode direction. In the case of a p-type sample, concentration of dominant Vþ is increased in the cathode direction by migration of positive charge vacancies in the reverse direction of the 3-field. Consequently, the growth rate in the cathode direction is accelerated due to the promotion of the adsorption of the Si atoms at the a-Si/NiSi2 and the Si atom rearrangement at the c-Si/NiSi2 phase boundary, which was found to be opposite compared to that for the n-type sample [34,43,44]. For doping time above 2 min, the MILC growth rate decreased toward a saturated value as shown in Fig. 2(d)-2(e), and, finally, it was saturated for the 5 min doping sample. Fig. 3(a)e3(c) show the FE-SEM images of morphology and microstructure of MILC and a-Si region corresponding to the 1 min B2H6 doped sample after 5 h annealing at 550 C, with the corresponding magnified FE-SEM images in the insets. The sample annealed without any 3-field is shown in Fig. 3(a), whereas those with applying external 3-field during annealing treatment in the cathode and the anode directions are shown in Fig. 3(b)e(c), respectively. The marks in Fig. 2(b)e(c) show the locations where the FE-SEM images in Fig. 3(a)-3(c) were taken. Preferred crystal growth orientations and the microstructures of the MILC growth region of 1min B2H6-doped sample with and without the applied 3-field in both the anode and the cathode directions are summarized in Table 1. The maximum crystallite lengths measured for an area of 10 mm 10 mm are given in Table 1 as well. The microstructure of the MILC region of the B2H6-doped sample without the 3-field showed a network of bi-directional
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needle like structures of branched crystallites with a regular angle of 70 to 110 between primary and secondary branches. The orientation of these branches are expected to be related with the orientation of NiSi2 precipitates in the a-Si film determining the subsequent lateral growth <111> direction of the crystallite [11,44]. The maximum length of branched crystallite was over 6.9 mm, with the width of 40e50 nm which differs from previous study results [37]. The microstructure of the B2H6-doped sample annealed in the anode direction of the 3-field showed an uni-directional needle like structure with preferred crystallite growth orientation of the <111> direction in the (110) plane [45]. Unlike samples grown without the 3-field (Fig. 3(a)), in the case of samples grown with the 3-field, the branching secondary arm is highly restrained, and a uniform array of parallel crystallites was formed as shown in Fig. 3(b)e(c). The length (over 11 mm) and the width (50e60 nm) of the crystallites with the 3-field at the leading edge of the MILC growth in the anode direction was 36% longer and 20% thicker, respectively, than those (6.9 mm and 40e50 nm) without the 3-field. The length (over 12 mm) and the width (60e80 nm) of the crystallites with the 3-field at the leading edge of the MILC growth in the cathode direction was 42% longer and 33% thicker, respectively, than those (6.9 mm and 40e50 nm) without the 3-field. This phenomenon may be attributed to the second level of Fig. 3(d) describing the different MILC growth rate of B2H6-dope a-Si sample in anode and cathode direction of 3-field along with the sample without the applied 3-field (explained in following paragraphs). The microstructure of the MILC region of B2H6-doped sample facing the cathode (Fig. 3(b)) also showed an uni-directional needle like structure. The maximum length of crystallite was over 12 mm with enhanced directionality compared to sample without electric field. This extension of crystallite length is associated with the long range migration of NiSi2 precipitates in the direction opposite the applied 3-field compared to the case without 3-field [40]. Meanwhile, the width of crystallite in the cathode direction was 60e80 nm, which was thicker than that (40e50 nm) of crystallite without the 3-field. In the case of PH3 doping [34], the MILC growth rate was decreased by increasing the level of doping, while in the case of B2H6 doping, the MILC growth rate increased compared to the PH3 doping and intrinsic Si. In the case of PH3-doped a-Si, the activated dopant (Pþ) suppresses the movement of the c-Si/NiSi2 interface through attraction of charged Ni ion which decelerate the MILC rate. However, in the case of B2H6 doped a-Si, the activated dopant (B) enhances the movement of the c-Si/NiSi2 interface through repulsion of charged Ni ion which accelerate the MILC rate. Fig. 3(d) shows the MILC growth rate of B2H6-doped a-Si thin films after 5 h annealing at 550 C in anode (þ, Blue) and cathode (-, Red) directions under an electric field (100 V/cm) together with the MILC growth rate without an electric field (Black) classified into three different MILC growth stages depending on the doping level. The MILC growth rate of the sample in anode direction was 14.54 mm/h in an intrinsic state and decreased to 11.82 mm/h after 1 min B2H6 doping. After further doping, the MILC growth rate decreased from 10.46 mm/h at 2 min to 10.24 mm/h at 4 min, and finally became saturated around 9.88 mm/h. On the other hand, without the 3-field, the growth rate increased from 7.38 mm/h to 10.08 mm/h when the doping level increased from no doping to 1 min doping. Above 1 min doping time, the growth rate became saturated around 10 mm/h. For the sample with the 3-field in the cathode direction, the growth rate was 5.78 mm/h for no doping and increased sharply to 14.42 mm/h after 1 min B2H6 doping. Above 1 min doping, the growth rate decreased to 11.74 mm/h after 2 min doping, and, finally, was saturated around 10.6 mm/h after 5 min doping. MILC growth rate of B2H6-doped a-Si thin films saturated regardless of the applied 3-field. In Fig. 3(d), the saturated MILC
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Fig. 3. (a) FE-SEM micrograph of the 1 min B2H6-doped a-Si thin films after the 550 C anneal for 5 h (no 3-field). FE-SEM micrographs of the 1 min B2H6-doped a-Si thin films in (b) the cathode and (c) the anode directions (under 3-field). (d) B2H6-doped a-Si thin film MILC growth rate for 0e5 min doping time in the anode (blue, þ) and the cathode (red, -) directions after the 550 C anneal for 5 h. MILC growth rate for the sample without 3-field is shown as black. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Table 1 Preferred Orientation and microstructure of MILC growth region for the sample doped with B2H6 for 1 min. Case
Dominant Microstructure
Expected Preferred Orientation
Maximum Length of Crystallite (mm)
Width of Needle Crystallite (nm)
Without ε-field With ε-field, anode
Uni-directional parallel growth Uni-directional parallel growth (with enhanced directionality) Uni-directional parallel growth (with enhanced directionality)
<111>/(110) <111>/(110)
>6.9 >11
40e50 50e60
<111>/(110)
>12
60e80
With ε-field, cathode
growth rate (~12 mm/hr) in the cathode direction was higher than that without an 3-field (~10 mm/hr), while that (~10 mm/hr) in the anode direction was similar. The saturated MILC growth rate (~10 mm/hr) in B2H6-doped a-Si was higher than that in intrinsic a-Si without an 3-field (7.3 mm/hr), while that (2 mm/hr) in PH3-doped aSi was lower than that in intrinsic a-Si without an 3-field (7.3 mm/hr) [34]. In the first part of Fig. 3(d), named “intrinsic stage,” the MILC growth rate of B2H6-doped samples is highest in the anode direction, intermediate without electric field, and lowest in the cathode direction. This order of MILC rate can be explained using the threestep hopping mechanism [38,43,44]. In the case of the anode direction, negatively charged Ni ions drift faster toward the leading interface edge of the anode direction, causing an enhancement of MILC growth rate. While, in the case of the cathode direction, due to the opposite direction of drift, MILC growth rate is decreased. However, in the absence of an 3-field, the MILC growth rate is elevated by the repulsive force between activated dopant (B) and negatively charged Ni ions. From Level I to Level II (“MILC rate reversal stage”), as doping
time is increased the MILC growth rate of the B2H6-doped sample in the cathode direction increased, whereas the decreasing trend in MILC growth rate was observed for the B2H6-doped sample in the anode direction. This phenomenon is associated with the positively-charged vacancy dependent MILC growth rate of B2H6doped a-Si thin film. The existence of positively-charged vacancies (Vþ and V2þ) in ptype Si can be used to explain the increasing and decreasing MILC growth rate in the cathodic and anodic directions, respectively. For p-type sample, as doping time is increased, the concentration of dominant positively-charged vacancies (Vþ and V2þ) is increased in the cathode direction due to their drifting toward the leading edge of NiSi2, which accelerates the MILC growth rate by increasing the adsorption and atomic rearrangement of the Si atoms. However, as doping time is increased more, the MILC growth rate in the cathode direction is slightly decreased due to establishment of an attractive force between activated dopant (B) and positively-charged vacancies. Decreasing the MILC growth rate as a function of doping time for p-type a-Si sample in the anode direction also can be understood with the same explanation. However, in this case,
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Fig. 4. Schematic diagrams of the MILC growth mechanisms for the B2H6-doped a-Si thin films under an applied 3-field. (a), (c), and (e) correspond to the anode direction, while (b), (d) and (f) corresponds to the cathode direction ((a) and (b) for doping level I, (c) and (d) for doping level II, (e) and (f) for doping level III).
concentration of positively charged vacancies is decreased in the anode direction. MILC growth rate is decelerated, and consequently the rate of the adsorption and the atomic rearrangement of Si atoms declines [46]. As seen from the third part of Fig. 3(d), the MILC growth rate is saturated after 3 min of B2H6 doping. This saturation behavior of MILC growth rate of B2H6-doped sample may be due to the internal interfacial electric field, similar to PH3-doped sample in our previous study [34]. Due to high doping time, the crystallization region of the Si thin film can be considered as having high concentration of activated stationary ions (B). On the other hand, a high density of states of defects in the mobility gap [47], holding holes and electrons, exists in the a-Si region, which may act as either p-type or n-type depending on the doping type. These defects can emit an extra electron and thus play a role as a donor. Therefore, it can be assumed that the a-Si region appears as a material with a donor state density (nx). Due to the existence of donor and acceptor atoms, a p-n junction is thought to be formed at the c-Si/a-Si interface. The internal 3-field at the interface can be estimated using similar equations as used for the n-type doping [34]. The estimated internal interfacial 3-field for the doping level of 1016 - 1020/cm3 was 3 104 V/cm to 3 MV/cm. The much higher internal interfacial 3-field induces the MILC growth rate saturation as the IMD doping time increases, for both with or with the applied external 3-field [34].
The operating mechanism and the rate determining step of pdoped MILC growth regions for different growth rate stages I, II, and III given in Fig. 4 are summarized in Table 2 as well. Filled and empty stars in Table 2 correspond to the factors enhancing and retarding the MILC, respectively: the greater the number of stars, the stronger the effect on the MILC growth rate. Doping Level I corresponds to the case of an intrinsic Si with or without 3-field, which is shown schematically in Fig. 4(a)e(b), where the rate of MILC is determined by the Ni ion and Ni vacancy hopping step. The case of doping Level II, which is a MILC rate reversal stage with charged vacancy migration affected by an external 3-field and slight p-type doping, is shown schematically in Fig. 4(c)e(d). The operating mechanism in doping Level I is still present, but the effect becomes minimal, which is indicated with one star. The other two major operating mechanisms, which are the repulsive interaction between B and Ni and the positively charged Si vacancies migration with the external electric field, are marked with three stars. The repulsive interaction between B and Ni enhances MILC both in the anode and the cathode direction, while the interaction between the positively charged Si vacancies with the electric field retards MILC in the anode direction and enhances it in the cathode direction. Since the two mechanisms are in the same direction in the cathode direction, the MILC rate was much enhanced in the cathode direction and resulted in the MILC rate reversal in the
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Table 2 Operating mechanism and rate determining step of boron-doped MILC growth regions for different growth stages I, II, and III. Mechanism
Doping Level I (Intrinsic Stage)
Without ε-field
① Nominal MILC by Ni ion and Ni vacancy hopping model @ 550 C
Doping Level II (MILC Rate Reversal Stage)
Doping Level III (MILC Rate Saturation Stage)
① MILC rate enhancement by B- and Ni- ① Saturation by large Internal interfacial ε-field due to repulsive interaction (+++) heavy doping (Internal interfacial ε-field » externally applied ε-field) (+++) ② MILC rate enhancement by B- and Ni- repulsive interaction (++) With ε-field - Anode ① MILC rate enhancement in the anode ① MILC rate enhancement by B- and Ni- ① Saturation by large Internal interfacial ε-field due to heavy doping (Internal interfacial ε-field » externally repulsive interaction (+++) Direction direction through Ni- movement ② Retardation of atomic rearrangement applied ε-field) (+++) enhancement by ε-field (+++) ② MILC rate enhancement by B- and Ni- repulsive by positively charged Si vacancy (Vþ) interaction (++) concentration decrease in the anode ③ Retardation of atomic rearrangement by positively direction (☆☆☆) ③ MILC rate enhancement in the anode charged Si vacancy (Vþ) concentration decrease in the direction through Ni- movement anode direction (☆) enhancement by ε-field (+) ④ MILC rate enhancement in the anode direction through Ni- movement enhancement by ε-field (+) With ε-field - Cathode ① MILC rate retardation in the cathode ① MILC rate enhancement by B- and Ni- ① Saturation by large Internal interfacial ε-field due to heavy doping (Internal interfacial ε-field » externally repulsive interaction (+++) Direction direction through Ni- movement applied ε-field) (+++) ② Enhancement of atomic retardation by ε-field (☆☆☆) rearrangement by positively charged Si ② MILC rate enhancement by B- and Ni- repulsive vacancy (Vþ) concentration increase in interaction (++) ③ Enhancement of atomic rearrangement by positively the cathode direction (+++) ③ MILC rate retardation in the cathode charged Si vacancy (Vþ) concentration increase in the cathode direction (+) direction through Ni- movement ④ MILC rate retardation in the cathode direction through retardation by ε-field (☆) Ni- movement retardation by ε-field (☆) Step 2 - Ni ion and Ni vacancy hopping Rate Determining Step Step 2 - Ni ion and Ni vacancy hopping Step 3 - Atomic rearrangement by positively charged Si vacancy (Vþ) migration
cathode direction. In the doping Level II, the rate determining step is the atomic rearrangement by positively-charged Si vacancy migration. Doping Level III corresponds to a MILC rate saturation stage with over doping effect in the a-Si region, as shown schematically in Fig. 4(e)e(f). The three operating mechanisms in doping Level II are still operating, but the effect is less significant. In the case of doping Level III, a fourth operating mechanism, which is an internal interfacial 3-field due to heavy doping, becomes the dominant operating mechanism. The internal interfacial 3-field takes over the other effects, which makes the MILC rate independent of 3-field direction. The internal interfacial 3-field becomes much greater than the externally applied 3-field in doping Level III, and the rate determining step is the Ni ion and Ni vacancy hopping step. The direction of the internal 3-field in the case of p-type doping is opposite to that of n-type doping, which causes the final saturated MILC rate to be higher in the case of p-type doping, and lower in the case of n-type doping compared to the intrinsic a-Si. In the case of p-type doping, the internal 3-field caused by the space charge region in between p-type c-Si and a-Si pushes the Ni ion away towards the a-Si region, while in the case of n-type doping, the internal 3-field pulls Ni towards the c-Si region. 4. Conclusions We have studied MILC growth rate saturation of heavily boron doped a-Si thin films by applying an external 3-field. The growth rate was classified into 3 stages based on the doping time. The operating mechanism and the rate determining step of p-doped MILC growth regions for three different growth stages are suggested. It was suggested that many competing driving forces are operating for p-type a-Si MILC growth, such as the repulsive force between the Ni ion and the activated B dopants, the forces between the Ni ions and the external 3-field, the positively charged vacancy migration, and the internal interfacial 3-field. For both the
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